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Cavity growth and filament formation of superplastically deformed Al 7475 Alloy
Materials Science and Engineering A298 (2001) 235 – 244 www.elsevier.com/locate/msea
Cavity growth and filament formation of superplastically deformed Al 7475 Alloy C.L. Chen *, M.J. Tan School of Mechanical and Production Engineering, Nanyang Technological Uni6ersity, Singapore 639798, Singapore Received 24 February 2000; received in revised form 1 June 2000
Abstract Superplastic deformation induced cavitation of aluminium alloys usually results in the material performance degradation. In this study the cavitation behaviour of Al 7475 was characterised, using samples deformed at temperature ranging from 480 to 530°C and at an initial strain rate of 10 − 3s − 1. The results showed that the cavity growth rate parameter increased slightly as the test temperature increased. The cavitation ratio at fracture increased evidently with increasing temperature from 480 to 500°C, and reached a plateau for any further increase of temperature. Various morphological filaments were observed at cavities and fracture surface, as evidence of the presence of liquid phase along grain boundaries. The effects of test temperature and thermal history on cavitation were found to be closely related to the presence of the liquid phase. The presence of liquid phase will improve the ability of the materials to tolerate high volume fraction of cavities before fracture. On the other hand, when liquid phase is anisotropically distributed along grain boundaries, it will cause the preferential interlinkage of cavities along the weak grain boundaries, and result in corporate grain boundary sliding (CGBS). It is concluded that the critical factor is to achieve appropriate quantity, high property, and uniformly distributed liquid phase along grain boundary. This highlights a new clue in searching for an economical and practical way to alleviate cavitation. © 2001 Elsevier Science B.V. All rights reserved. Keywords: Superplasticity; Filament formation; Cavitation; Thermal history
1. Introduction During the past two decades, much effort has been devoted to the development of superplastic forming (SPF) high strength aluminium alloys and forming techniques, since high strength aluminium alloys have wide applications in aerospace industry. Unfortunately, cavitation usually occurs in a wide range of metallic materials and metal matrix composites (MMCs) after SPF, especially in aluminium alloys, leading to the degradation of the overall properties of the post-SPF materials [1,2]. It has been demonstrated that the mechanical properties of the materials would be significantly reduced when the cavity ratio exceeded approximately 1% [3]. This limits the useable strain range for SPF to a certain value at which the overall properties of the materials are not seriously degraded by cavitation, though many SPF materials can be * Corresponding author. Tel.: +65-7906122; fax: +65-7906122. E-mail address: [email protected] (C.L. Chen).
strained to as high as 1000% or even more before fracture. In order to alleviate the problem of cavitation, studies of the effects of superimposed pressure, hot isostatic presses (HIP), and prior-deformation heat treatments, etc., have been extensively carried out recently [4–7]. The superimposed pressure and HIP methods need relatively high additional costs, and besides, or more importantly, the dimension of the forming components is limited using these techniques. From the practical point of view, prior-deformation heat treatment is an economic and applicable way to alleviate cavitation by altering the thermal history of the materials [5]. The thermal history plays an important role in the cavitation behaviour of the materials. The variation of cavitation behaviour for samples with different thermal history is usually accompanied by the changes of filament formation [8,9]. The filament formation of superplastic deformed materials is frequently reported. It is generally thought as the evidence of the existence of liquid phase during the deformation [8–11]. Cao et
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al. [8] suggested that the filament formation was more likely to be a viscous flow of a liquid-like substance originating at the grain boundaries. Blandin et al. [9] found that filaments present on the deformed samples of as-received materials but the filaments are no longer present on heat-treated ones. The presence of liquid phase was expected to greatly change the properties of grain boundaries, and to influence the cavitation behaviour of the materials. The altering of thermal history through heat treatment could be a potential way to modify the role of liquid phase. However, little information is available on these aspects. The research therefore aimed to investigate the nature of liquid phase and its effect on cavity growth. It is also expected that the results of this research would be useful for the understanding of cavitation mechanism and the searching of economic and applicable methods to alleviate cavitation. 2. Experimental procedures Commercial Al 7475 sheet of thickness 2.0 mm used in the present experiment has a chemical composition of (wt.%) 5.64 Zn, 2.34 Mg, 1.58 Cu, 0.19 Cr, 0.08 Fe, 0.05 Si, 0.02 Ti, 0.01 Mn and Al balance. The as-received sheet was in T4 condition (solution heat treated, quenched and natural aged). The tensile specimens were of 15 mm length, 4 mm width and 2 mm thickness in gauge parts, with sample cutting direction parallel to the rolling direction of the sheet. The surfaces of tensile sample gauge part were polished up to 1 mm diamond paste.
The optimum superplastic deformation condition of Al 7475 has been reported at strain rate around 10 − 3 s − 1, and at temperature about 516°C [9,12–14]. In the present work, samples were uniaxial deformed at temperature ranging from 480°C to 530°C and at an initial strain rate of 10 − 3 s − 1, in a furnace chamber in air at atmospheric pressure. The time to raise the temperature to desired value was about 1 h. After the temperature reached the desired value, stabilisation for 15 min was used for all samples before the tensile test start. To investigate the influence of thermal history, some samples were first heated up to 516°C or 530°C, and held for 15 min; after that, temperature was lowered to 480°C within 5 min and held for another 5 min to homogenise temperature; finally, samples were strained to fracture at 480°C. Some samples experienced additional 20% strain at 516°C or 530°C before the lowering of the test temperature. Following straining to a desired strain or to fracture, samples were unloaded immediately and were cooled down to room temperature using force cooling, instead of quenching in water, to preserve clean surface for scanning electron microscope (SEM) observation. For cavitation ratio measurement, samples were sectioned, mounted, ground and polished up to 1 mm diamond paste. Cavitation ratios were determined on longitudinal-short traverse (L-N) section of samples using image analyser. Samples for microstructure observation were etched with Keller’s reagent. SEM was used for surface morphology and microstructure observation and energy dispersive X-ray spectroscopy (EDS) analysis was employed to determine the composition of filaments on sample surface.
3. Results and discussion
3.1. As recei6ed condition of commercial Al 7475
Fig. 1. Microstructure of as received Al 7475 sheet.
The microstructure of as-received Al 7475 sheet showed that the material was fully recrystallised, Fig. 1. The grains were in pancake-shape, i.e. equiaxed shape in longitudinal-transverse (L-T) section and thin in short transverse (N) direction. Its average linear intercepts were 10 mm, 10.1 mm and 7.2 mm in the L, T and N directions, respectively. Precipitates could be found in the interior grain and along grain boundaries. The precipitates were generally about 0.5 mm or less in diameter. EDS analysis showed that the precipitates were rich in Al, Mg, Zn, and Cu. Some Fe and Cr rich precipitates were also detected. Mg, Zn, Fe and Cu enriched precipitates were frequently reported for the Al 7475 [9,10,13,15].
C.L. Chen, M.J. Tan / Materials Science and Engineering A298 (2001) 235–244 Table 1 List of elongation to fracture, of (%), at initial strain rate of 10−3s−1 Specimen code
Test procedure
of (%)
AR1
480°C, 10−3 s−1, strained to fracture 500°C, 10−3 s−1, strained to fracture 516°C, 10−3 s−1, strained to fracture 530°C, 10−3 s−1, strained to fracture 530°C, 15 min; 480°C, 10−3 s−1, strained to fracture 516°C, 15 min; 480°C, 10−3 s−1, strained to fracture 530°C, 15 min; 530°C, 10−3 s−1, strained to 20%; 480°C, 10−3 s−1, strained to fracture 516°C, 15min; 516°C, 10−3 s−1, strained to 20%; 480°C, 10−3 s−1, strained to fracture
395 910%
AR2 AR3 AR4 HT1 HT2 HT3
HT4
690 910% 850 910% 730 910% 510915% 720910% 320 915%
400915%
Fig. 2. Strain – stress curves of different thermal history samples strained at 480°C.
3.2. Superplastic deformation and ca6ity growth 3.2.1. Elongation to failure The elongation to fracture of samples are listed in Table 1. The elongation to fracture of as-received samples tested at temperature range from 480°C to 530°C peaked at 516°C, viz. 850%. At 480°C, the elongation was quite low. However, the sample configuration after fracture revealed uniform deformation at all temperature used. The present results were in good agreement with the optimum superplastic deformation condition of Al 7475 reported in the previous investigations [9,12–14]. The elongation of samples with different thermal history tested at 480°C are listed together in Table 1. Their strain–stress curves are showed in Fig. 2. It is interesting to note in Table 1 and Fig. 2 that, after
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exposure at 516°C for 15 min, the elongation to fracture of sample strained later at 480°C was nearly doubled. 530°C exposure also showed beneficial effects though the improvement is relatively small. The beneficial effect of high temperature exposure may be attributed to the partial shrinkage of the pre-existing cavities and dissolution of grain boundary precipitates [5,9]. The decrease of sm in Fig. 2 of the sample after 530°C or 516°C exposure implied that the grain boundaries might be smoother due to the partial dissolution of grain boundary precipitates. Hence, larger effects of the precipitate dissolution could be expected for the higher temperature exposure. This inference was in consistent to the fact that the reduction of sm after 530°C exposure was larger than that of 516°C exposure. The reason behind the 530°C exposure had less beneficial effect on elongation than the 516°C exposure needs further clarification. The precipitate dissolution and cavity shrinkage at the 530°C exposure should be more significant, which predicates contradict conclusions to the results in the Table 1. This can not be well explained by the grain growth during high temperature exposure, since microstructure examination showed limited grain growth during exposure. Furthermore, the serious decrease of the elongation to fracture at 480°C after 20% strain at 530 or 516°C is difficult to be explained by these factors. Some other factors related to the different thermal histories may account for it. This will be discussed in Section 3.5.2.
3.2.2. Ca6ity growth It is well accepted that grain boundary sliding (GBS) is the dominate mechanism of superplasticity, and that sliding needs to be accommodated by other mechanisms, for example, grain-boundary migration, recrystallisation, diffusion flow, or dislocation slip. When GBS cannot be well accommodated, stress concentration at certain sites may cause the development of cavitation. The triple junction of grains, particles and even ledges on grain boundary can serve as the nucleate sites for cavities [16–18]. The stable nuclei or pre-existing cavity then grows through stress-directed vacancy diffusion and plastic deformation of the surrounding materials. With the increase of cavity radius, for example, to above 1 mm, the contribution from diffusion is quickly reduced and the plasticity controlled growth is dominant [19,20]. There are some theoretical work concerned with plasticity controlled void growth during superplastic flow [21–25]. They all predicted a similar relationship initially proposed by Hancock [21] of the type dV/dt = KVo; , however, their predictions of K value vary by considering the physical situation of a void. According to Hancock’s model [21], taking K as a constant, the cavity ratio under plasticity controlled growth can be expressed by Cv = C0 exp(ho), where h (identical to K)
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Fig. 3. Variation of cavitation ratio with strain of samples after deformation at high temperature. Label of 516–480°C refer to the samples experienced 15 min exposure at 516°C and then strained at 480°C.
Table 2 List of cavity growth parameters Temperature
C0
h
Cf
of
480°C 500°C 516°C 530°C 516°C–480°C
0.0023 0.0014 0.0003 0.0002 0.0003
2.07 2.41 2.97 3.33 2.61
0.060 0.201 0.194 0.196 0.068
3959 10% 690 9 10% 850910% 730 9 10% 720 9 10%
is the cavity growth rate parameter, C0 is a constant, o is the true strain. Cavity ratio variations with strain for the present materials are shown in Fig. 3 Cavity growth parameters after regression to Hancock’s model are listed in Table 2. The cavity ratios at fracture, Cf,
measured along gauge part with the final fracture region excluded, are also listed in Table 2.Cf reflects the cavity level the material can withstand, beyond which unstable traverse cavity interlink may take place and result in the final fracture. Three aspects were noted in Fig. 3 and Table 2. First, C0 declined as the test temperature was raised from 480 to 530°C. This lowering of C0 can result from the greater beneficial effects of the closing up of the pre-existing cavities during heating up and 15 min stabilisation at higher temperature. Secondly, the materials deformed above 500°C had a relatively high Cf, of about 0.2; while sample tested at 480°C had a relatively low Cf, of about 0.06. Finally, h increased slightly with the increase of test temperature. The later two aspects are related to the effects of the liquid phase. The relatively advantageous cavitation behaviour of the sample deformed at 516°C can be interpreted by its appropriate combination of C0 and h; while the longest elongation is also attributed to the high Cf. On the other hand, the lowest Cf account for the lowest elongation of materials deformed at 480°C. The microstructure showed that samples deformed at 480°C usually has a large amount of small cavities, Fig. 4(a). Optical microscope observation showed that the cavity amount tend to increase with the increase of strain. This can result from either the continuous nucleation [5,7] of new cavities or the small cavities that were initially indistinguishable grew to visible size. Hosokawa et al. [26] had reported a similar result on an Al-Mg alloy. For samples tested at 530°C, cavities were usually larger in dimension but less in quantity, Fig. 4(b). Compared with the case of 480°C, the less quantity of cavity can result from: (1) the decrease of continuous nucleation, (2) the increase of cavity coalescence, (3) both (1) and (2), and more probably, (4) the cavity coalescence outweighs the continuous nucleation of cavity. Therefore, it is suggested that, at 530°C, plasticity controlled cavity growth and interlink of cavities
Fig. 4. Cavity morphology of sample strained to 300% at (a) 480; and (b) 530°C. Stressed horizontally.
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Fig. 5. Morphology of filaments of samples tested at (a) 480; (b) 500; (c) 516; and (d) 530°C. Stressed horizontally.
along longitudinal direction are more important and result in the large size cavity, as well as large h and Cf values.
3.3. Filament formation 3.3.1. Morphology of filaments Formation of filament on Al 7475 after superplastic deformation has been previously reported [8 –11]. The filament formation is generally thought to be the evidence of the existence of liquid phase along grain boundaries at high temperature. The temperature of filament formation was reported to be as low as 437°C [8]. The Al-Zn-Mg system has phases of melting points at 437 and 382°C [27,28]. The liquid phase can be resulted from the dissolution of incipient melting phases distributed adjacent to grain boundaries at high temperature [29]. Filaments were formed from the liquid phase when the grain boundaries were separated during superplastic deformation. As the filaments were formed from the liquid phase, the morphology and composition of the filament should be closely related to the nature of the liquid phase. Notably, in the present investigation, the morphology
of the filament changed as the deformation temperature changed, as shown in Fig. 5. At 480°C, filaments were short and fine, of about 0.2 mm in diameter, Fig. 5(a). The amount of filaments on samples deformed at 480°C was less than those tested at higher temperature. By increasing test temperature up to 516°C, filaments became denser, longer and coarser. The filament diameters were about 0.8 mm and 1.2 mm for samples tested at 500 and 516°C, respectively, Fig. 5(b) and (c). At 530°C, thin and wide laminate-liked filaments developed. Some were above 5 mm in width, Fig. 5(d). Long and continuous filaments, some above 120 mm in length, more than 10 times of the grain diameter, were observed on narrow and long surface cavities of samples deformed at 516°C, as marked by points A and B in Fig. 6(a). The occurrence of such long filaments declined in the following test temperature order: 516, 500 and 530°C. Long filament was not observed in the case of 480°C, as shown in Fig. 6(b). The continuous filaments linking disconnected grains indicated the relative movement of grains such as GBS. This implied that GBS could be very large in some regions and its effect on the cavity development was significant.
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3.3.2. EDS analysis The EDS spectrums are shown in Fig. 7. As can be seen from Fig. 7(a), filaments formed at 480°C consisted of Al, small amount of Mg, Zn and even smaller quantity of Cu. In the case of 516°C, high Mg and Zn enrichment and slight Cu enrichment could be detected on filaments. Enrichment of Mg, Zn and Cu may be related to the dissolution of incipient melting phase or precipitates along grain boundaries at high temperature. As the temperature was raised from 480 to 516°C, dissolution of more precipitates resulted in the enrichment of Mg, Zn and Cu in liquid phase and then in filaments. The EDS spectrums detected along filament were nearly the same on samples deformed at 480 or 500°C. For samples tested at 516 or 530°C, variations of the composition along filament were noted. Fig. 7(b) presents the typical EDS spectrums detected at different positions along filament of the sample tested at 516°C. Relatively high Mg and Zn peaks were detected at the center part (point C) and intermediate part (point B) of the filament shown in Fig. 5(c), compared with the bottom part (point A) and inside the grain (point D). At 530°C, the enrichment of Mg and Zn was high at the filament center and it reduced nearly to the level of inside grain at the filament bottom. If the bottom part of filaments represents the later-formed liquid and the center part represents the earlier-formed liquid, the composition variations indicated that the Mg and Zn rich precipitates dissolved during deformation and precipitates were nearly exhausted in the later deformation stage, especially for the case of 530°C. For samples exposed at 516 or 530°C for 15 min and then deformed at 480°C, Mg and Zn enrichment along filament was quite uniform, Fig. 7(c). This indicated that the precipitates dissolved slowly and were not exhausted in this test condition.
3.4. Thermodynamics analysis The formation of filament was difficult to be interpreted by either dislocation or diffusion mechanism. Cao et al. [8] suggested that the filament was formed through viscous flow of a liquid-like or semi-liquid grain boundary materials. With the liquid phase present along grain boundaries, the viscous flow can be expressed as o; = tm, where m is the viscosity, t is the stress. In this case, the viscosity of the liquid will play an important role during deformation. For most liquid metals the variation of viscosity m with temperature T(K) can be written as
m = m0 exp
E , RT
(1)
where m0 and E are constants, R is the gas constant. For liquid Al, Mg and Zn, m0 and E are 0.1492, 0.0245, 0.4131 mPa · s and 16.5, 30.5, 12.7 kJ/ mol [30,31], respectively. At the present test temperature range of 753–803 K, extrapolations according to Eq. (1) showed that the viscosity of Mg and Zn are larger than that of Al. The high enrichment of Mg and Zn in liquid phase therefore was expected to increase the viscosity of liquid phase. In addition to the viscosity variation when the liquid composition changed, the surface free energy of the liquid phase would influence the filament formation, since the filament had high surface volume ratio. The surface free energy is the surface tension plus the free energy due to the surface concentrations of the components, i.e. Gs = s+% G( i Gi,
(2)
where Gs is the surface free energy, s is the surface tension, G( i is the partial molar free energy of the component i, and Gi is the surface concentration of component i. At stable state, Gs is at minimum. At
Fig. 6. Surface cavity morphology of samples tested at (a) 516; (b) 480°C. Stressed horizontally.
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Korolkov [32] reported that surface tension of liquid aluminium decreased when the Mg and Zn contents increased. The reduction is about 30% for the presence of 5–6 wt.% Mg and relatively small for Zn. Thus, these solutes, especially Mg, would concentrate in the filament surface. The concentration of Mg in filament surface explains the concurrent Mg and O peaks in EDS spectrums in Fig. 7. Fox [33] reported that when the concentration of Mg is greater than 0.2wt.%, MgO will form preferentially to Al2O3. Since the Pilling-Bedworth ratio of Mg is 0.78, which is less than 1, the oxide film formed has less protection effect from further oxidation. Serious oxidation then takes place.
3.5. Effect of liquid phase
Fig. 7. EDS spectrums detected (a) at filament bottom of samples deformed at 480 – 530°C. (b) Along filament of samples deformed at 516°C (Points A, B. C and D were corresponded to Fig. 5c. (c) On filament of samples with different thermal history deformed at 480°C.
constant temperature, according to Gibbs adsorption isotherm, there is Gi = −
1 ds , RT d ln ai
(3)
where ai is the activity of constituent i in liquid. It indicates that any constituent which lowers s has a positive value of Gi, i.e. it will concentrate on the surface.
3.5.1. Ca6ity nucleation and growth The presence of liquid along grain boundaries significantly changed the cavitation behaviour of the material. The presence of liquid phase could relax stress concentration at junctions of grain boundaries or particles, which is beneficial for the uniform deformation of the materials. On the other hand, liquid phase would reduce the grain boundary strength [34], which may lead to the acceleration of the cavity nucleation and growth. That the material can sustain high cavitation ratio at fracture when strained above 500°C may relate to the presence of adequate amount of liquid phase at grain boundaries. Adequate amount of liquid leads to the quick relaxation of stress concentration and uniform deformation of the material. Thus, the material can withstand high cavity volume fraction before fracture. At 480°C, melting of incipient phase at grain boundaries was relatively small. As a result, the effect of liquid phase is limited and the material can tolerate only a small volume fraction of cavity. The nucleation of cavity under stress was related to the energy needed for the newly created surface. With the liquid phase presented along grain boundaries, the nucleation of cavity would depend on the surface energy of the liquid, Gs According to Raj and Ashby [35], the radius of a stable cavity under stress t is rc =2Gs/t. Generally, Gs increases with decreasing temperature [36]. Simultaneously, the flow stress, t, also increased when the temperature decreased. Comparing these two aspects, the increase of stress may outweigh the increase of surface energy when temperature decreases. As a result, the stable cavity radius decreases with the decrease in test temperature. Thus, more cavities would nucleate when temperature decreased, as reflected by cavity morphology changes described in Section 3.2.2. The viscosity of the liquid is an important factor which affects the grain boundary strength [34,37]. The higher viscosity liquid phase would bind the grain boundaries strongly. Since enrichment of Mg and Zn are expected to increase the viscosity of the liquid
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phase, the increased of Mg, Zn contents in liquid hence alleviated the decrease of the viscosity of the liquid when the temperature was raised. This means that the grain boundary strength decreased slowly as the temperature increased. The variation of grain boundary strength explains the slight increase of cavity growth rate parameter when the deformation temperature was increased from 480 to 530°C.
3.5.2. Effect of thermal history The influence of thermal history on cavitation was related to the liquid property variation. For samples after exposure at 516 or 530°C, many pre-existing cavities were closed up and precipitates were dissolved. The liquid phase along grain boundary then contained high Mg and Zn, which increased liquid viscosity. After cooling down to 480°C, grain boundaries contained more liquid than boundaries of samples without exposure. This was evidenced by the denser and longer filaments formed during later straining, Fig. 8. As discussed above, reasonable increase in liquid quantity could be beneficial for stress concentration relaxation while increase of liquid viscosity could lower the cavity
Fig. 8. Filament morphology of sample after exposed at 530°C for 15 min and strained to fracture at 480°C.
growth rate parameter. The high temperature exposure thus would lower C0 result in h value between the cases of 480 and 516°C, and make samples deformed more uniformly, as shown in Table 2 and Fig. 9(b) Consequently, the cavitation was alleviated and greater elongation was achieved. The anisotropic distribution of liquid phase along grain boundaries accounted for the preferentially longitudinal interlinkage of cavities at high temperature. Since the rolling process usually leads to a break-up and an alignment of the second phase particles along the rolling direction [16,18,38], cavities formed at these particles were lined up along the rolling direction at earlier straining. At 530 and 516°C, continuous dissolution of precipitates during straining caused the grain boundaries parallel to stress direction covering with more liquid, though the earlier formed liquid was gradually squeezed to the boundaries perpendicular to stress direction. Those grain boundaries close to longitudinal direction therefore had lower bonding strength. The cavities thus interlink easily along these boundaries, as shown in Fig. 10. For samples deformed at 480°C, only a small quantity of liquid was formed at grain boundaries, and interlinkage of cavities along longitudinal direction was limited. Under the applied stress, traverse interlinkage of cavities and corporate grain boundary sliding (CGBS), which were frequently observed, resulted in the final fracture, Fig. 9(a). The deleterious effect of 20% strain at 516 or 530°C before straining at 480°C could be related to the redistribution of liquid phase along grain boundaries during the first period of straining, as illustrated in Fig. 11. At 516 or 530°C, the low viscous liquid was easily squeezed out of grain boundaries parallel to the stress direction and transferred to perpendicular ones [37]. This caused those grain boundaries parallel to the stress direction to contain less liquid phase at high temperature after straining. After being cooled down to 480°C, these grain boundaries would be lack of liquid and hence had greater bound strength than those boundaries perpendicular to stress direction where more liquid existed, since the decrease in liquid quantity causes the rapid increase of the viscosity of the viscous solid phase at grain boundary [37]. Anisotropic grain boundary property distribution resulted in CGBS, as shown in Fig. 9(c), because sliding along some grain boundaries parallel to the stress direction was difficult. Serious interlinkage of cavities took place along the boundaries in traverse direction, leading to premature failure.
3.6. Summary
Fig. 10. Preferentially longitudinal interlinkage of cavities at 516°C, o= 640%, stressed horizontally.
Cavitation behaviour is closely related to the presence of the liquid phase. The materials with appropriate quantity of liquid phase along grain boundaries can
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Fig. 9. Surface morphology of samples after (a) strained to fracture at 480°C; (b) 516°C exposure for 15 min; strained to fracture at 480°C; (c) 530°C exposure for 15 min and strained to 20%; strained to fracture at 480°C. Stressed horizontally.
tolerate high volume fraction of cavities. On the other hand, the liquid phase causes the grain boundary strength to decrease and result in larger cavity growth rate. The critical factor, as have been illustrated in the present investigation, is to achieve appropriate quantity, high property, and uniformly distributed liquid phase along grain boundary. Heat treatment therefore could be an economic and practical method to control the cavity behaviour.
The variations of the liquid quantity and viscosity account for the slight increase of cavity growth rate with increasing temperature. Short time exposure at temperature higher than forming temperature of the materials before SPF is suggested while the deformation during the short time higher temperature exposure should be avoided.
4. Conclusions Filaments were observed at cavities and fracture surface as evidence of liquid phase along grain boundaries. The quantity and radius of the filament increased as testing temperature increases. Continuous filaments emerged on long and narrow cavities indicating that GBS was closely related to cavity growth. The deformation temperature and thermal history significantly influence the filament quantity and composition. The presence of Mg and Zn enrichment is expected to improve the viscosity of the liquid phase which bound the grain boundaries. Cavity grows preferentially along the weak grain boundaries, explaining the cavity interlink behaviour. CGBS may take place when anisotropic grain boundary property distribution is involved.
Fig. 11. Illustration of liquid distribution at different deformation stage.
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References [1] R.D. Schelleng, G.H. Reynolds, Metall. Trans. 4 (1973) 2199. [2] N. Ridley, D.W. Livesey, A.K. Mukherjee, Metall Trans. 15A (A) (1984) 1443. [3] Ph. Bompard, J.Y. Lacroix, A. Varloteaux, Aluminium 64 (1988) 162. [4] J. Pilling, N. Ridley, Acta Metall. 34 (1986) 669. [5] A. Varloteaux, J.J. Blandin, M. Suery, Mater. Sci. Tech. 5 (1989) 1109. [6] Xun Yuwei, Zhun Yiyuan, Mao Wenfeng, Cui Jianzhong, J. Mater. Proc. Tech. 72 (1997) 183–187. [7] C.C. Bampton, J.W. Edington, Metall. Trans. 13A (A, October) (1982) 1721 – 1727. [8] W.D. Cao, X.P. Lu, H. Conrad, Acta mater. 44 (2) (1996) 697 – 706. [9] J.J. Brandin, B. Hong, A. Varloteaux, M. Suery, G. L’esperanve, Acta Mater. 44 (1996) 2317–2326. [10] Y Takayama, T. Tozawa, H. Kato, Acta mater. 47 (4) (1999) 1263 – 1270. [11] M.G. Zelin, Acta mater. 45 (9) (1997) 3533–3542. [12] John A. Wert, 1982. J. of Metals, Sept., 35–41. [13] Dong Hyuk Shin, Chong Soo Lee, Woo-Jin Kim, Acta mater. 45 (12) (1997) 5195 –5202. [14] R.K. Mahidhara, J. Mat. Eng. and Perf. 4 (6, December) (1995) 674 – 678. [15] H. Yoshida, Mater. Sci. Forum 204–206 (1996) 657. [16] A.H. Chokshi, Mater. Sci. Forum 233–234 (1997) 89 – 108. [17] H. Iwasaki, M. Takeuchi, T. Mori, M. Mabuchi, K. Higashi, Scripta Metallurgica et, Materialia 31 (3) (1994) 255– 260. [18] Dong. H. Shin, Kyung-Tae Park, Materials Science and Engineering A268 (1999) 55–62. [19] N. Ridley, Z.C. Wang, Mat. Sci. Forun. 233–234 (1997) 63 – 80]. [20] N. Ridley, Z.C. Wang, Mat. Sci. Forun. 170–172 (1994) 177 – 186.
.
[21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33]
[34] [35] [36]
[37] [38]
J.W. Hancock, Meatl. Sci. 10 (1976) 319 – 325. A. Needleman, J.R. Rice, Acta Metall. 28 (1980) 1315–1332. A.C.F. Cocks, M.F. Ashby, Met. Sci. 14 (1980) 395–402. M.J. Stowell, D.W. Livesey, N. Ridley, Acta Metall. 32 (1984) 35 – 42. D.S. Wilkinson, C.H. Caceres, Mat. Sci. Tech. 2 (1986) 1086– 1092. H. Hosokawa, H. Iwasaki, T. Mori, M. Mabuchi, T. Tagata, K. Higashi, Acta Mater. 47 (6) (1999) 1859 – 1867. J.B. Clark, F.N. Rhines, Trans. Am. Inst. Min. Engrs. 209 (1957) 6. G.R. Goldak, J.G. Parr, J. Inst. Met. 92 (1963/4) 230. M.C. Dang, J.J. Blandin, B. Baudelet, Acta Mater. 44 (10) (1996) 3991 – 4002. F. King, Aluminium and Its Alloys, Market Cross House, Cooper Street, Chichester, West Sussex, England, 1987. Eric A. Brandes, Smithells Metals Reference Book, sixth edition, Butterworth & Co. (Publishers) Ltd, 1983, pp. 14 – 17. A.M. Korolkov, Casting Properties of Metals and Alloys, Consultants Bureau, New York, 1960. Fox, M.H., Nilmani, M., 1993. Aluminium cast house technology, theory and practice. In: Milmani, M., (Ed.), Aluminium Melt Treatment and Casting, 4 – 8, July, 1993, University of Melbourne, The Minerals, Metals and Materials Society, Australia. W.J. Kim, J. Wolfenstine, O.D. Sherby, Acta Metall. Mater. 39 (1991) 199 – 208. R. Raj, M.F. Ashby, Acta metall. 23 (1975) 653 – 666. Yogeshwar Sahai, 1993. Aluminium cast house technology, theory and practice. In: Milmani, M., (Ed.), Aluminium Melt Treatment and Casting, 4 – 8, July, 1993, University of Melbourne, The Minerals, Metals and Materials Society, Australia, pp 265 – 279. M.C. Dang, B. Baudelet, J. Mat. Sci. 29 (1994) 2315–2320. H. Iwasaki, T. Mori, M. Mabuchi, K. Higashi, Mater. Sci. Technol. 15 (1999) 180 – 184.
Cavity growth and filament formation of superplastically deformed Al 7475 Alloy C.L. Chen *, M.J. Tan School of Mechanical and Production Engineering, Nanyang Technological Uni6ersity, Singapore 639798, Singapore Received 24 February 2000; received in revised form 1 June 2000
Abstract Superplastic deformation induced cavitation of aluminium alloys usually results in the material performance degradation. In this study the cavitation behaviour of Al 7475 was characterised, using samples deformed at temperature ranging from 480 to 530°C and at an initial strain rate of 10 − 3s − 1. The results showed that the cavity growth rate parameter increased slightly as the test temperature increased. The cavitation ratio at fracture increased evidently with increasing temperature from 480 to 500°C, and reached a plateau for any further increase of temperature. Various morphological filaments were observed at cavities and fracture surface, as evidence of the presence of liquid phase along grain boundaries. The effects of test temperature and thermal history on cavitation were found to be closely related to the presence of the liquid phase. The presence of liquid phase will improve the ability of the materials to tolerate high volume fraction of cavities before fracture. On the other hand, when liquid phase is anisotropically distributed along grain boundaries, it will cause the preferential interlinkage of cavities along the weak grain boundaries, and result in corporate grain boundary sliding (CGBS). It is concluded that the critical factor is to achieve appropriate quantity, high property, and uniformly distributed liquid phase along grain boundary. This highlights a new clue in searching for an economical and practical way to alleviate cavitation. © 2001 Elsevier Science B.V. All rights reserved. Keywords: Superplasticity; Filament formation; Cavitation; Thermal history
1. Introduction During the past two decades, much effort has been devoted to the development of superplastic forming (SPF) high strength aluminium alloys and forming techniques, since high strength aluminium alloys have wide applications in aerospace industry. Unfortunately, cavitation usually occurs in a wide range of metallic materials and metal matrix composites (MMCs) after SPF, especially in aluminium alloys, leading to the degradation of the overall properties of the post-SPF materials [1,2]. It has been demonstrated that the mechanical properties of the materials would be significantly reduced when the cavity ratio exceeded approximately 1% [3]. This limits the useable strain range for SPF to a certain value at which the overall properties of the materials are not seriously degraded by cavitation, though many SPF materials can be * Corresponding author. Tel.: +65-7906122; fax: +65-7906122. E-mail address: [email protected] (C.L. Chen).
strained to as high as 1000% or even more before fracture. In order to alleviate the problem of cavitation, studies of the effects of superimposed pressure, hot isostatic presses (HIP), and prior-deformation heat treatments, etc., have been extensively carried out recently [4–7]. The superimposed pressure and HIP methods need relatively high additional costs, and besides, or more importantly, the dimension of the forming components is limited using these techniques. From the practical point of view, prior-deformation heat treatment is an economic and applicable way to alleviate cavitation by altering the thermal history of the materials [5]. The thermal history plays an important role in the cavitation behaviour of the materials. The variation of cavitation behaviour for samples with different thermal history is usually accompanied by the changes of filament formation [8,9]. The filament formation of superplastic deformed materials is frequently reported. It is generally thought as the evidence of the existence of liquid phase during the deformation [8–11]. Cao et
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al. [8] suggested that the filament formation was more likely to be a viscous flow of a liquid-like substance originating at the grain boundaries. Blandin et al. [9] found that filaments present on the deformed samples of as-received materials but the filaments are no longer present on heat-treated ones. The presence of liquid phase was expected to greatly change the properties of grain boundaries, and to influence the cavitation behaviour of the materials. The altering of thermal history through heat treatment could be a potential way to modify the role of liquid phase. However, little information is available on these aspects. The research therefore aimed to investigate the nature of liquid phase and its effect on cavity growth. It is also expected that the results of this research would be useful for the understanding of cavitation mechanism and the searching of economic and applicable methods to alleviate cavitation. 2. Experimental procedures Commercial Al 7475 sheet of thickness 2.0 mm used in the present experiment has a chemical composition of (wt.%) 5.64 Zn, 2.34 Mg, 1.58 Cu, 0.19 Cr, 0.08 Fe, 0.05 Si, 0.02 Ti, 0.01 Mn and Al balance. The as-received sheet was in T4 condition (solution heat treated, quenched and natural aged). The tensile specimens were of 15 mm length, 4 mm width and 2 mm thickness in gauge parts, with sample cutting direction parallel to the rolling direction of the sheet. The surfaces of tensile sample gauge part were polished up to 1 mm diamond paste.
The optimum superplastic deformation condition of Al 7475 has been reported at strain rate around 10 − 3 s − 1, and at temperature about 516°C [9,12–14]. In the present work, samples were uniaxial deformed at temperature ranging from 480°C to 530°C and at an initial strain rate of 10 − 3 s − 1, in a furnace chamber in air at atmospheric pressure. The time to raise the temperature to desired value was about 1 h. After the temperature reached the desired value, stabilisation for 15 min was used for all samples before the tensile test start. To investigate the influence of thermal history, some samples were first heated up to 516°C or 530°C, and held for 15 min; after that, temperature was lowered to 480°C within 5 min and held for another 5 min to homogenise temperature; finally, samples were strained to fracture at 480°C. Some samples experienced additional 20% strain at 516°C or 530°C before the lowering of the test temperature. Following straining to a desired strain or to fracture, samples were unloaded immediately and were cooled down to room temperature using force cooling, instead of quenching in water, to preserve clean surface for scanning electron microscope (SEM) observation. For cavitation ratio measurement, samples were sectioned, mounted, ground and polished up to 1 mm diamond paste. Cavitation ratios were determined on longitudinal-short traverse (L-N) section of samples using image analyser. Samples for microstructure observation were etched with Keller’s reagent. SEM was used for surface morphology and microstructure observation and energy dispersive X-ray spectroscopy (EDS) analysis was employed to determine the composition of filaments on sample surface.
3. Results and discussion
3.1. As recei6ed condition of commercial Al 7475
Fig. 1. Microstructure of as received Al 7475 sheet.
The microstructure of as-received Al 7475 sheet showed that the material was fully recrystallised, Fig. 1. The grains were in pancake-shape, i.e. equiaxed shape in longitudinal-transverse (L-T) section and thin in short transverse (N) direction. Its average linear intercepts were 10 mm, 10.1 mm and 7.2 mm in the L, T and N directions, respectively. Precipitates could be found in the interior grain and along grain boundaries. The precipitates were generally about 0.5 mm or less in diameter. EDS analysis showed that the precipitates were rich in Al, Mg, Zn, and Cu. Some Fe and Cr rich precipitates were also detected. Mg, Zn, Fe and Cu enriched precipitates were frequently reported for the Al 7475 [9,10,13,15].
C.L. Chen, M.J. Tan / Materials Science and Engineering A298 (2001) 235–244 Table 1 List of elongation to fracture, of (%), at initial strain rate of 10−3s−1 Specimen code
Test procedure
of (%)
AR1
480°C, 10−3 s−1, strained to fracture 500°C, 10−3 s−1, strained to fracture 516°C, 10−3 s−1, strained to fracture 530°C, 10−3 s−1, strained to fracture 530°C, 15 min; 480°C, 10−3 s−1, strained to fracture 516°C, 15 min; 480°C, 10−3 s−1, strained to fracture 530°C, 15 min; 530°C, 10−3 s−1, strained to 20%; 480°C, 10−3 s−1, strained to fracture 516°C, 15min; 516°C, 10−3 s−1, strained to 20%; 480°C, 10−3 s−1, strained to fracture
395 910%
AR2 AR3 AR4 HT1 HT2 HT3
HT4
690 910% 850 910% 730 910% 510915% 720910% 320 915%
400915%
Fig. 2. Strain – stress curves of different thermal history samples strained at 480°C.
3.2. Superplastic deformation and ca6ity growth 3.2.1. Elongation to failure The elongation to fracture of samples are listed in Table 1. The elongation to fracture of as-received samples tested at temperature range from 480°C to 530°C peaked at 516°C, viz. 850%. At 480°C, the elongation was quite low. However, the sample configuration after fracture revealed uniform deformation at all temperature used. The present results were in good agreement with the optimum superplastic deformation condition of Al 7475 reported in the previous investigations [9,12–14]. The elongation of samples with different thermal history tested at 480°C are listed together in Table 1. Their strain–stress curves are showed in Fig. 2. It is interesting to note in Table 1 and Fig. 2 that, after
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exposure at 516°C for 15 min, the elongation to fracture of sample strained later at 480°C was nearly doubled. 530°C exposure also showed beneficial effects though the improvement is relatively small. The beneficial effect of high temperature exposure may be attributed to the partial shrinkage of the pre-existing cavities and dissolution of grain boundary precipitates [5,9]. The decrease of sm in Fig. 2 of the sample after 530°C or 516°C exposure implied that the grain boundaries might be smoother due to the partial dissolution of grain boundary precipitates. Hence, larger effects of the precipitate dissolution could be expected for the higher temperature exposure. This inference was in consistent to the fact that the reduction of sm after 530°C exposure was larger than that of 516°C exposure. The reason behind the 530°C exposure had less beneficial effect on elongation than the 516°C exposure needs further clarification. The precipitate dissolution and cavity shrinkage at the 530°C exposure should be more significant, which predicates contradict conclusions to the results in the Table 1. This can not be well explained by the grain growth during high temperature exposure, since microstructure examination showed limited grain growth during exposure. Furthermore, the serious decrease of the elongation to fracture at 480°C after 20% strain at 530 or 516°C is difficult to be explained by these factors. Some other factors related to the different thermal histories may account for it. This will be discussed in Section 3.5.2.
3.2.2. Ca6ity growth It is well accepted that grain boundary sliding (GBS) is the dominate mechanism of superplasticity, and that sliding needs to be accommodated by other mechanisms, for example, grain-boundary migration, recrystallisation, diffusion flow, or dislocation slip. When GBS cannot be well accommodated, stress concentration at certain sites may cause the development of cavitation. The triple junction of grains, particles and even ledges on grain boundary can serve as the nucleate sites for cavities [16–18]. The stable nuclei or pre-existing cavity then grows through stress-directed vacancy diffusion and plastic deformation of the surrounding materials. With the increase of cavity radius, for example, to above 1 mm, the contribution from diffusion is quickly reduced and the plasticity controlled growth is dominant [19,20]. There are some theoretical work concerned with plasticity controlled void growth during superplastic flow [21–25]. They all predicted a similar relationship initially proposed by Hancock [21] of the type dV/dt = KVo; , however, their predictions of K value vary by considering the physical situation of a void. According to Hancock’s model [21], taking K as a constant, the cavity ratio under plasticity controlled growth can be expressed by Cv = C0 exp(ho), where h (identical to K)
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Fig. 3. Variation of cavitation ratio with strain of samples after deformation at high temperature. Label of 516–480°C refer to the samples experienced 15 min exposure at 516°C and then strained at 480°C.
Table 2 List of cavity growth parameters Temperature
C0
h
Cf
of
480°C 500°C 516°C 530°C 516°C–480°C
0.0023 0.0014 0.0003 0.0002 0.0003
2.07 2.41 2.97 3.33 2.61
0.060 0.201 0.194 0.196 0.068
3959 10% 690 9 10% 850910% 730 9 10% 720 9 10%
is the cavity growth rate parameter, C0 is a constant, o is the true strain. Cavity ratio variations with strain for the present materials are shown in Fig. 3 Cavity growth parameters after regression to Hancock’s model are listed in Table 2. The cavity ratios at fracture, Cf,
measured along gauge part with the final fracture region excluded, are also listed in Table 2.Cf reflects the cavity level the material can withstand, beyond which unstable traverse cavity interlink may take place and result in the final fracture. Three aspects were noted in Fig. 3 and Table 2. First, C0 declined as the test temperature was raised from 480 to 530°C. This lowering of C0 can result from the greater beneficial effects of the closing up of the pre-existing cavities during heating up and 15 min stabilisation at higher temperature. Secondly, the materials deformed above 500°C had a relatively high Cf, of about 0.2; while sample tested at 480°C had a relatively low Cf, of about 0.06. Finally, h increased slightly with the increase of test temperature. The later two aspects are related to the effects of the liquid phase. The relatively advantageous cavitation behaviour of the sample deformed at 516°C can be interpreted by its appropriate combination of C0 and h; while the longest elongation is also attributed to the high Cf. On the other hand, the lowest Cf account for the lowest elongation of materials deformed at 480°C. The microstructure showed that samples deformed at 480°C usually has a large amount of small cavities, Fig. 4(a). Optical microscope observation showed that the cavity amount tend to increase with the increase of strain. This can result from either the continuous nucleation [5,7] of new cavities or the small cavities that were initially indistinguishable grew to visible size. Hosokawa et al. [26] had reported a similar result on an Al-Mg alloy. For samples tested at 530°C, cavities were usually larger in dimension but less in quantity, Fig. 4(b). Compared with the case of 480°C, the less quantity of cavity can result from: (1) the decrease of continuous nucleation, (2) the increase of cavity coalescence, (3) both (1) and (2), and more probably, (4) the cavity coalescence outweighs the continuous nucleation of cavity. Therefore, it is suggested that, at 530°C, plasticity controlled cavity growth and interlink of cavities
Fig. 4. Cavity morphology of sample strained to 300% at (a) 480; and (b) 530°C. Stressed horizontally.
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Fig. 5. Morphology of filaments of samples tested at (a) 480; (b) 500; (c) 516; and (d) 530°C. Stressed horizontally.
along longitudinal direction are more important and result in the large size cavity, as well as large h and Cf values.
3.3. Filament formation 3.3.1. Morphology of filaments Formation of filament on Al 7475 after superplastic deformation has been previously reported [8 –11]. The filament formation is generally thought to be the evidence of the existence of liquid phase along grain boundaries at high temperature. The temperature of filament formation was reported to be as low as 437°C [8]. The Al-Zn-Mg system has phases of melting points at 437 and 382°C [27,28]. The liquid phase can be resulted from the dissolution of incipient melting phases distributed adjacent to grain boundaries at high temperature [29]. Filaments were formed from the liquid phase when the grain boundaries were separated during superplastic deformation. As the filaments were formed from the liquid phase, the morphology and composition of the filament should be closely related to the nature of the liquid phase. Notably, in the present investigation, the morphology
of the filament changed as the deformation temperature changed, as shown in Fig. 5. At 480°C, filaments were short and fine, of about 0.2 mm in diameter, Fig. 5(a). The amount of filaments on samples deformed at 480°C was less than those tested at higher temperature. By increasing test temperature up to 516°C, filaments became denser, longer and coarser. The filament diameters were about 0.8 mm and 1.2 mm for samples tested at 500 and 516°C, respectively, Fig. 5(b) and (c). At 530°C, thin and wide laminate-liked filaments developed. Some were above 5 mm in width, Fig. 5(d). Long and continuous filaments, some above 120 mm in length, more than 10 times of the grain diameter, were observed on narrow and long surface cavities of samples deformed at 516°C, as marked by points A and B in Fig. 6(a). The occurrence of such long filaments declined in the following test temperature order: 516, 500 and 530°C. Long filament was not observed in the case of 480°C, as shown in Fig. 6(b). The continuous filaments linking disconnected grains indicated the relative movement of grains such as GBS. This implied that GBS could be very large in some regions and its effect on the cavity development was significant.
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3.3.2. EDS analysis The EDS spectrums are shown in Fig. 7. As can be seen from Fig. 7(a), filaments formed at 480°C consisted of Al, small amount of Mg, Zn and even smaller quantity of Cu. In the case of 516°C, high Mg and Zn enrichment and slight Cu enrichment could be detected on filaments. Enrichment of Mg, Zn and Cu may be related to the dissolution of incipient melting phase or precipitates along grain boundaries at high temperature. As the temperature was raised from 480 to 516°C, dissolution of more precipitates resulted in the enrichment of Mg, Zn and Cu in liquid phase and then in filaments. The EDS spectrums detected along filament were nearly the same on samples deformed at 480 or 500°C. For samples tested at 516 or 530°C, variations of the composition along filament were noted. Fig. 7(b) presents the typical EDS spectrums detected at different positions along filament of the sample tested at 516°C. Relatively high Mg and Zn peaks were detected at the center part (point C) and intermediate part (point B) of the filament shown in Fig. 5(c), compared with the bottom part (point A) and inside the grain (point D). At 530°C, the enrichment of Mg and Zn was high at the filament center and it reduced nearly to the level of inside grain at the filament bottom. If the bottom part of filaments represents the later-formed liquid and the center part represents the earlier-formed liquid, the composition variations indicated that the Mg and Zn rich precipitates dissolved during deformation and precipitates were nearly exhausted in the later deformation stage, especially for the case of 530°C. For samples exposed at 516 or 530°C for 15 min and then deformed at 480°C, Mg and Zn enrichment along filament was quite uniform, Fig. 7(c). This indicated that the precipitates dissolved slowly and were not exhausted in this test condition.
3.4. Thermodynamics analysis The formation of filament was difficult to be interpreted by either dislocation or diffusion mechanism. Cao et al. [8] suggested that the filament was formed through viscous flow of a liquid-like or semi-liquid grain boundary materials. With the liquid phase present along grain boundaries, the viscous flow can be expressed as o; = tm, where m is the viscosity, t is the stress. In this case, the viscosity of the liquid will play an important role during deformation. For most liquid metals the variation of viscosity m with temperature T(K) can be written as
m = m0 exp
E , RT
(1)
where m0 and E are constants, R is the gas constant. For liquid Al, Mg and Zn, m0 and E are 0.1492, 0.0245, 0.4131 mPa · s and 16.5, 30.5, 12.7 kJ/ mol [30,31], respectively. At the present test temperature range of 753–803 K, extrapolations according to Eq. (1) showed that the viscosity of Mg and Zn are larger than that of Al. The high enrichment of Mg and Zn in liquid phase therefore was expected to increase the viscosity of liquid phase. In addition to the viscosity variation when the liquid composition changed, the surface free energy of the liquid phase would influence the filament formation, since the filament had high surface volume ratio. The surface free energy is the surface tension plus the free energy due to the surface concentrations of the components, i.e. Gs = s+% G( i Gi,
(2)
where Gs is the surface free energy, s is the surface tension, G( i is the partial molar free energy of the component i, and Gi is the surface concentration of component i. At stable state, Gs is at minimum. At
Fig. 6. Surface cavity morphology of samples tested at (a) 516; (b) 480°C. Stressed horizontally.
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Korolkov [32] reported that surface tension of liquid aluminium decreased when the Mg and Zn contents increased. The reduction is about 30% for the presence of 5–6 wt.% Mg and relatively small for Zn. Thus, these solutes, especially Mg, would concentrate in the filament surface. The concentration of Mg in filament surface explains the concurrent Mg and O peaks in EDS spectrums in Fig. 7. Fox [33] reported that when the concentration of Mg is greater than 0.2wt.%, MgO will form preferentially to Al2O3. Since the Pilling-Bedworth ratio of Mg is 0.78, which is less than 1, the oxide film formed has less protection effect from further oxidation. Serious oxidation then takes place.
3.5. Effect of liquid phase
Fig. 7. EDS spectrums detected (a) at filament bottom of samples deformed at 480 – 530°C. (b) Along filament of samples deformed at 516°C (Points A, B. C and D were corresponded to Fig. 5c. (c) On filament of samples with different thermal history deformed at 480°C.
constant temperature, according to Gibbs adsorption isotherm, there is Gi = −
1 ds , RT d ln ai
(3)
where ai is the activity of constituent i in liquid. It indicates that any constituent which lowers s has a positive value of Gi, i.e. it will concentrate on the surface.
3.5.1. Ca6ity nucleation and growth The presence of liquid along grain boundaries significantly changed the cavitation behaviour of the material. The presence of liquid phase could relax stress concentration at junctions of grain boundaries or particles, which is beneficial for the uniform deformation of the materials. On the other hand, liquid phase would reduce the grain boundary strength [34], which may lead to the acceleration of the cavity nucleation and growth. That the material can sustain high cavitation ratio at fracture when strained above 500°C may relate to the presence of adequate amount of liquid phase at grain boundaries. Adequate amount of liquid leads to the quick relaxation of stress concentration and uniform deformation of the material. Thus, the material can withstand high cavity volume fraction before fracture. At 480°C, melting of incipient phase at grain boundaries was relatively small. As a result, the effect of liquid phase is limited and the material can tolerate only a small volume fraction of cavity. The nucleation of cavity under stress was related to the energy needed for the newly created surface. With the liquid phase presented along grain boundaries, the nucleation of cavity would depend on the surface energy of the liquid, Gs According to Raj and Ashby [35], the radius of a stable cavity under stress t is rc =2Gs/t. Generally, Gs increases with decreasing temperature [36]. Simultaneously, the flow stress, t, also increased when the temperature decreased. Comparing these two aspects, the increase of stress may outweigh the increase of surface energy when temperature decreases. As a result, the stable cavity radius decreases with the decrease in test temperature. Thus, more cavities would nucleate when temperature decreased, as reflected by cavity morphology changes described in Section 3.2.2. The viscosity of the liquid is an important factor which affects the grain boundary strength [34,37]. The higher viscosity liquid phase would bind the grain boundaries strongly. Since enrichment of Mg and Zn are expected to increase the viscosity of the liquid
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phase, the increased of Mg, Zn contents in liquid hence alleviated the decrease of the viscosity of the liquid when the temperature was raised. This means that the grain boundary strength decreased slowly as the temperature increased. The variation of grain boundary strength explains the slight increase of cavity growth rate parameter when the deformation temperature was increased from 480 to 530°C.
3.5.2. Effect of thermal history The influence of thermal history on cavitation was related to the liquid property variation. For samples after exposure at 516 or 530°C, many pre-existing cavities were closed up and precipitates were dissolved. The liquid phase along grain boundary then contained high Mg and Zn, which increased liquid viscosity. After cooling down to 480°C, grain boundaries contained more liquid than boundaries of samples without exposure. This was evidenced by the denser and longer filaments formed during later straining, Fig. 8. As discussed above, reasonable increase in liquid quantity could be beneficial for stress concentration relaxation while increase of liquid viscosity could lower the cavity
Fig. 8. Filament morphology of sample after exposed at 530°C for 15 min and strained to fracture at 480°C.
growth rate parameter. The high temperature exposure thus would lower C0 result in h value between the cases of 480 and 516°C, and make samples deformed more uniformly, as shown in Table 2 and Fig. 9(b) Consequently, the cavitation was alleviated and greater elongation was achieved. The anisotropic distribution of liquid phase along grain boundaries accounted for the preferentially longitudinal interlinkage of cavities at high temperature. Since the rolling process usually leads to a break-up and an alignment of the second phase particles along the rolling direction [16,18,38], cavities formed at these particles were lined up along the rolling direction at earlier straining. At 530 and 516°C, continuous dissolution of precipitates during straining caused the grain boundaries parallel to stress direction covering with more liquid, though the earlier formed liquid was gradually squeezed to the boundaries perpendicular to stress direction. Those grain boundaries close to longitudinal direction therefore had lower bonding strength. The cavities thus interlink easily along these boundaries, as shown in Fig. 10. For samples deformed at 480°C, only a small quantity of liquid was formed at grain boundaries, and interlinkage of cavities along longitudinal direction was limited. Under the applied stress, traverse interlinkage of cavities and corporate grain boundary sliding (CGBS), which were frequently observed, resulted in the final fracture, Fig. 9(a). The deleterious effect of 20% strain at 516 or 530°C before straining at 480°C could be related to the redistribution of liquid phase along grain boundaries during the first period of straining, as illustrated in Fig. 11. At 516 or 530°C, the low viscous liquid was easily squeezed out of grain boundaries parallel to the stress direction and transferred to perpendicular ones [37]. This caused those grain boundaries parallel to the stress direction to contain less liquid phase at high temperature after straining. After being cooled down to 480°C, these grain boundaries would be lack of liquid and hence had greater bound strength than those boundaries perpendicular to stress direction where more liquid existed, since the decrease in liquid quantity causes the rapid increase of the viscosity of the viscous solid phase at grain boundary [37]. Anisotropic grain boundary property distribution resulted in CGBS, as shown in Fig. 9(c), because sliding along some grain boundaries parallel to the stress direction was difficult. Serious interlinkage of cavities took place along the boundaries in traverse direction, leading to premature failure.
3.6. Summary
Fig. 10. Preferentially longitudinal interlinkage of cavities at 516°C, o= 640%, stressed horizontally.
Cavitation behaviour is closely related to the presence of the liquid phase. The materials with appropriate quantity of liquid phase along grain boundaries can
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Fig. 9. Surface morphology of samples after (a) strained to fracture at 480°C; (b) 516°C exposure for 15 min; strained to fracture at 480°C; (c) 530°C exposure for 15 min and strained to 20%; strained to fracture at 480°C. Stressed horizontally.
tolerate high volume fraction of cavities. On the other hand, the liquid phase causes the grain boundary strength to decrease and result in larger cavity growth rate. The critical factor, as have been illustrated in the present investigation, is to achieve appropriate quantity, high property, and uniformly distributed liquid phase along grain boundary. Heat treatment therefore could be an economic and practical method to control the cavity behaviour.
The variations of the liquid quantity and viscosity account for the slight increase of cavity growth rate with increasing temperature. Short time exposure at temperature higher than forming temperature of the materials before SPF is suggested while the deformation during the short time higher temperature exposure should be avoided.
4. Conclusions Filaments were observed at cavities and fracture surface as evidence of liquid phase along grain boundaries. The quantity and radius of the filament increased as testing temperature increases. Continuous filaments emerged on long and narrow cavities indicating that GBS was closely related to cavity growth. The deformation temperature and thermal history significantly influence the filament quantity and composition. The presence of Mg and Zn enrichment is expected to improve the viscosity of the liquid phase which bound the grain boundaries. Cavity grows preferentially along the weak grain boundaries, explaining the cavity interlink behaviour. CGBS may take place when anisotropic grain boundary property distribution is involved.
Fig. 11. Illustration of liquid distribution at different deformation stage.
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References [1] R.D. Schelleng, G.H. Reynolds, Metall. Trans. 4 (1973) 2199. [2] N. Ridley, D.W. Livesey, A.K. Mukherjee, Metall Trans. 15A (A) (1984) 1443. [3] Ph. Bompard, J.Y. Lacroix, A. Varloteaux, Aluminium 64 (1988) 162. [4] J. Pilling, N. Ridley, Acta Metall. 34 (1986) 669. [5] A. Varloteaux, J.J. Blandin, M. Suery, Mater. Sci. Tech. 5 (1989) 1109. [6] Xun Yuwei, Zhun Yiyuan, Mao Wenfeng, Cui Jianzhong, J. Mater. Proc. Tech. 72 (1997) 183–187. [7] C.C. Bampton, J.W. Edington, Metall. Trans. 13A (A, October) (1982) 1721 – 1727. [8] W.D. Cao, X.P. Lu, H. Conrad, Acta mater. 44 (2) (1996) 697 – 706. [9] J.J. Brandin, B. Hong, A. Varloteaux, M. Suery, G. L’esperanve, Acta Mater. 44 (1996) 2317–2326. [10] Y Takayama, T. Tozawa, H. Kato, Acta mater. 47 (4) (1999) 1263 – 1270. [11] M.G. Zelin, Acta mater. 45 (9) (1997) 3533–3542. [12] John A. Wert, 1982. J. of Metals, Sept., 35–41. [13] Dong Hyuk Shin, Chong Soo Lee, Woo-Jin Kim, Acta mater. 45 (12) (1997) 5195 –5202. [14] R.K. Mahidhara, J. Mat. Eng. and Perf. 4 (6, December) (1995) 674 – 678. [15] H. Yoshida, Mater. Sci. Forum 204–206 (1996) 657. [16] A.H. Chokshi, Mater. Sci. Forum 233–234 (1997) 89 – 108. [17] H. Iwasaki, M. Takeuchi, T. Mori, M. Mabuchi, K. Higashi, Scripta Metallurgica et, Materialia 31 (3) (1994) 255– 260. [18] Dong. H. Shin, Kyung-Tae Park, Materials Science and Engineering A268 (1999) 55–62. [19] N. Ridley, Z.C. Wang, Mat. Sci. Forun. 233–234 (1997) 63 – 80]. [20] N. Ridley, Z.C. Wang, Mat. Sci. Forun. 170–172 (1994) 177 – 186.
.
[21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33]
[34] [35] [36]
[37] [38]
J.W. Hancock, Meatl. Sci. 10 (1976) 319 – 325. A. Needleman, J.R. Rice, Acta Metall. 28 (1980) 1315–1332. A.C.F. Cocks, M.F. Ashby, Met. Sci. 14 (1980) 395–402. M.J. Stowell, D.W. Livesey, N. Ridley, Acta Metall. 32 (1984) 35 – 42. D.S. Wilkinson, C.H. Caceres, Mat. Sci. Tech. 2 (1986) 1086– 1092. H. Hosokawa, H. Iwasaki, T. Mori, M. Mabuchi, T. Tagata, K. Higashi, Acta Mater. 47 (6) (1999) 1859 – 1867. J.B. Clark, F.N. Rhines, Trans. Am. Inst. Min. Engrs. 209 (1957) 6. G.R. Goldak, J.G. Parr, J. Inst. Met. 92 (1963/4) 230. M.C. Dang, J.J. Blandin, B. Baudelet, Acta Mater. 44 (10) (1996) 3991 – 4002. F. King, Aluminium and Its Alloys, Market Cross House, Cooper Street, Chichester, West Sussex, England, 1987. Eric A. Brandes, Smithells Metals Reference Book, sixth edition, Butterworth & Co. (Publishers) Ltd, 1983, pp. 14 – 17. A.M. Korolkov, Casting Properties of Metals and Alloys, Consultants Bureau, New York, 1960. Fox, M.H., Nilmani, M., 1993. Aluminium cast house technology, theory and practice. In: Milmani, M., (Ed.), Aluminium Melt Treatment and Casting, 4 – 8, July, 1993, University of Melbourne, The Minerals, Metals and Materials Society, Australia. W.J. Kim, J. Wolfenstine, O.D. Sherby, Acta Metall. Mater. 39 (1991) 199 – 208. R. Raj, M.F. Ashby, Acta metall. 23 (1975) 653 – 666. Yogeshwar Sahai, 1993. Aluminium cast house technology, theory and practice. In: Milmani, M., (Ed.), Aluminium Melt Treatment and Casting, 4 – 8, July, 1993, University of Melbourne, The Minerals, Metals and Materials Society, Australia, pp 265 – 279. M.C. Dang, B. Baudelet, J. Mat. Sci. 29 (1994) 2315–2320. H. Iwasaki, T. Mori, M. Mabuchi, K. Higashi, Mater. Sci. Technol. 15 (1999) 180 – 184.